Interface complex reinforced Fe-Co-M based martensitic alloy and preparation method thereof

Abstract

The invention discloses an interface complex reinforced Fe-Co-M-based martensitic alloy and a preparation method thereof. A microscopic phase of the Fe-Co-M-based martensitic alloy comprises a small-angle grain boundary structure modified by a complex composed of a high-density dislocation structure and metal-interstitial atoms. The element percentage composition expression is FeaCobMcQd, a is greater than or equal to 40% and less than or equal to 60%, b is greater than or equal to 30% and less than or equal to 50%, c is greater than 0 and less than or equal to 14%, d is greater than 0 and less than or equal to 2, and a + b + c + d is equal to 100; m comprises at least one of Mo, V, Ti, W, Nb, Ta, Cr, Mn, Zr, Hf, Si, P, Ga, In, Sn, Pb, Ge, As, Sb, Te, Cu, Zn, Au, Ag, Pt, Pd, Cd, Ru, K, Mg, Y and Ca; q comprises at least one of B, C and N. According to the martensitic alloy, a small-angle interface martensitic structure modified by a high-density interface complex is obtained through pre-deformation machining and a low-temperature annealing process, so that the synergistic strengthening effect of interface strengthening, solid solution strengthening and dislocation strengthening of a nanoscale large-angle grain boundary is achieved, and the martensitic alloy has the ultrahigh tensile yield strength of a 3 GPa-level block material; and meanwhile, 5% tensile plasticity is kept.

Description

Technical Field

[0001] The present invention relates to a martensitic alloy, in particular to an interface complex-reinforced Fe-Co-M-based martensitic alloy and a preparation method thereof, belonging to the technical field of metal materials. Background Art

[0002] Ultra-high-strength bulk metallic structural materials occupy an irreplaceable position in advanced structural applications such as aerospace, transportation, and energy. With the continuous advancement of modern industry, the demand for ultra-high-strength metallic materials to achieve strength goals while maintaining usable ductility has become increasingly urgent. Alloys characterized by a lath martensite structure have long been the epitome of ultra-high-strength alloys. Lath martensite is the product of a diffusionless austenite-to-martensite transformation during rapid cooling of martensitic alloys. A notable feature of this material is its microstructure, which contains an ultra-dense dislocation network and finely distributed martensitic interfaces. This can enhance the yield strength of the single-phase BCC martensitic matrix to 1–1.4 GPa. Notably, these martensitic interfaces are composed of a high proportion (50–70 vol.%) of lath and sub-bulk grain boundaries, primarily low-angle grain boundaries. These low-angle grain boundaries strengthen the matrix primarily through dislocation strengthening, accounting for over 65% of the total yield strength. However, although these low-angle grain boundaries have good coherence with the matrix, which is beneficial to promote dislocation transfer at the interface to accommodate plastic strain, they also fundamentally damage their ability to hinder dislocation motion, thereby limiting the strength of martensitic alloys.

[0003] To further enhance the strength of martensitic alloys, conventional alloy design approaches typically involve the introduction of high-density precipitation or hierarchical multiphase substructures. These approaches have met with limited success, typically capping their strength at around 2 GPa. The reasons for these limitations are multiple: precipitation strengthening is limited by equilibrium solubility, microstructure-related dislocation strengthening is limited by thermodynamic recovery, and the yield strength of the hierarchical multiphase substructure is limited by the soft phase or soft zone. These inherent limitations, coupled with the inherent permeability of small-angle grain boundaries in martensite, have resulted in a persistent trade-off between strength and ductility. The yield strength of the most advanced ultra-high-strength alloys reported to date rarely exceeds 2.5 GPa, and their corresponding tensile ductility is typically less than 3%.

[0004] The aforementioned methods for regulating the mechanical properties of martensitic alloys are essentially based on increasing the defect density in the alloy to strengthen the alloy by introducing a second phase, high-density dislocations, or hierarchical multiphase substructures. Although increasing the defect density in martensitic alloys can improve the strength of martensitic alloys to a certain extent, the inherent limitations of these methods and the limited ability of martensitic lath interfaces to block dislocations have hindered further breakthroughs in the strength of martensitic alloys. These inherent limitations include:

[0005] 1. The introduction of precipitation strengthening is limited by equilibrium solubility, resulting in an upper limit on its volume fraction.

[0006] 2. Dislocation strengthening in the hierarchical multiphase substructure is limited by thermodynamic recovery, and the hierarchical multiphase structure is limited by the soft phase or soft area, resulting in limited improvement in yield strength.

[0007] 3. The high density and low angle grain boundaries in lath martensite are prone to dislocation penetration, which limits the strength improvement. Summary of the Invention

[0008] To address the problems encountered in prior art where precipitation strengthening is limited by equilibrium solubility, resulting in an upper limit on its volume fraction and thus limiting strength gains; dislocation strengthening in hierarchical multiphase substructures is limited by thermodynamic recovery; and the hierarchical multiphase microstructure is limited by soft phases or soft zones, resulting in limited yield strength gains; and the inherently permeable nature of high-density low-angle grain boundaries in lath martensite limits strength gains. The present invention provides an interface-complex-strengthened Fe-Co-M-based martensitic alloy, characterized by a high-density dislocation structure and low-angle grain boundaries modified by metal-interstitial atom complexes. This alloy exhibits both excellent deformation resistance and plastic deformation, achieving a comprehensive toughness that surpasses the strength of martensitic alloys while maintaining usable plasticity. This Fe-Co-M-based martensitic alloy introduces an interface complex (a chemical complex composed of metal and interstitial atoms) through interface engineering. This alloy enhances its dislocation-blocking effect by modulating the mechanical stability of defects, thereby improving the alloy's strength. This approach differs from the traditional approach of increasing alloy strength by increasing density.

[0009] A second object of the present invention is to provide a method for preparing an interface complex-strengthened Fe-Co-M-based martensitic alloy. This method, by regulating the content and type of alloying elements and combining cold working deformation with heat treatment, obtains a high-density dislocation structure and a low-angle grain boundary structure modified by a complex composed of metal-interstitial atoms in the Fe-Co-M-based martensitic alloy, thereby achieving its strengthening and toughening. The method has the advantages of simple operation and short process flow, and can be produced in batches and on a large scale.

[0010] In order to achieve the above technical objectives, the present invention provides an interface composite strengthened Fe-Co-M based martensitic alloy, the element percentage composition expression of which is: Fe a Co b M c Q d, wherein, 40%≤a≤60%, 30%≤b≤50%, 0<c≤14%, 0<d≤2%, a+b+c+d =100; M includes at least one of Mo, V, Ti, W, Nb, Ta, Cr, Mn, Zr, Hf, Si, P, Ga, In, Sn, Pb, Ge, As, Sb, Te, Cu, Zn, Au, Ag, Pt, Pd, Cd, Ru, K, Mg, Y, and Ca; Q includes at least one of B, C, and N; its microphase comprises a high-density dislocation structure and a low-angle grain boundary structure modified by a complex composed of metal-interstitial atoms.

[0011] The key to the Fe-Co-M based martensitic alloy of the present invention is that it adopts the interface complex method to strengthen the alloy strength, which can surpass the limits of existing conventional alloy strength strengthening methods. The present invention adopts the interface complex strategy to promote small-angle grain boundaries to hinder dislocation migration to the greatest extent, while suppressing crack nucleation, while traditional multiphase strengthening methods are limited by equilibrium solubility. The present invention utilizes high-density dislocations and their derived boundaries as defect regulators, and develops a Fe-Co-M medium-entropy martensitic alloy system with component combination optimization and appropriate thermomechanical processing design, wherein the Co element can stabilize the body-centered cubic (BCC) martensitic structure, while the M metal element has a large mismatch between the atomic size and the matrix element Fe or Co, resulting in severe lattice distortion, providing energy and chemical driving force for subsequent element segregation. The martensitic matrix has a high defect density (dislocation density > 10 -15 m -2 ), which produces a synergistic strengthening effect with lattice distortion and solid solution strengthening. Meanwhile, the high-density defects during annealing provide fast pathways for element diffusion. Furthermore, the introduction of interstitial atoms, such as B (0.3 at.%) and C (0.1 at.%), improves hardenability and enhances interfacial bonding.

[0012] The present invention further prefers Fe-Co-M based martensitic alloy, whose element percentage composition expression is: Fe a Co b M c Q d , among which, 40%≤a≤60%, 30%≤b≤50%, 1<c≤12%, 0<d≤2%, a+b+c+d =100.

[0013] The present invention also provides a method for preparing an interface complex-strengthened Fe-Co-M-based martensitic alloy, which comprises the following steps: pre-treating the surfaces of the metal raw material M, Fe, and Co, and pre-treating the non-metallic raw material Q by crushing; adding the pre-treated metal raw material and non-metallic raw material into a smelting furnace; and sequentially performing vacuum melting, casting, homogenization, cold deformation, and annealing to obtain the alloy.

[0014] Based on the introduction of M metal elements and interstitial atoms into the Fe-Co alloy, the present invention can further optimize the alloy properties through appropriate thermal processing technology. For example, cold rolling increases the dislocation density, and subsequent low-temperature annealing triggers interface segregation to promote the formation of interface complexes, stabilize defects, and ultimately optimize the mechanical properties of the alloy.

[0015] As a preferred solution, the metal raw materials are ultrasonically cleaned with alcohol. The present invention can remove organic pollutants on the surface of the raw materials Fe, Co, and M metal elements through ethanol ultrasonic cleaning, thereby reducing the introduction of impurity elements into the alloy material.

[0016] As a preferred solution, the non-metallic raw materials are crushed to a particle size of -1 mm. Crushing the Q non-metallic raw materials to an appropriate particle size is beneficial to promoting the mutual reaction between the various raw materials.

[0017] As a preferred solution, during the vacuum melting process, metal raw materials are stacked in the melting furnace from top to bottom in order of melting point, while non-metallic raw materials are placed at the very bottom. By optimizing the placement order of the metal and non-metallic raw materials, the high-melting-point metal raw materials can be fully melted while miniscule interstitial atoms are protected from being blown away by the flame generated by the tungsten electrode.

[0018] As a preferred embodiment, the vacuum melting is achieved in an arc melting furnace. The arc melting conditions are: an arc current of 370-390 A, a melting time of 90-120 s per melting, and a cooling period of 4-5 minutes after each melting. The sample is then flipped and the next melting is repeated 5-8 times. The present invention achieves a uniform melt through thorough melting.

[0019] As a preferred solution, the homogenization treatment conditions are: in an inert gas atmosphere, at a temperature of 1373-1523 K, for 1.5-2 hours. The inert atmosphere is provided by an inert gas, such as argon, with a pressure of 0.4-0.6 atmospheres. Before filling the inert gas, the furnace chamber needs to be evacuated to 2×10 -1 Pa, and in order to minimize the O2 content in the container, the vacuum-filling-inert gas steps will be repeated 3 to 5 times.

[0020] As a preferred solution, the cold deformation treatment conditions are: rolling at room temperature, a total rolling deformation of 70-90%, a pass deformation of 3-5%, and a roller speed of 4-8 r / min. Under the preferred cold deformation treatment conditions, the sample dimensions should be processed with a smaller downward pressure but multiple passes, ultimately obtaining a plate rolling sample with a larger rolling volume, which can ensure that more energy is stored in the sample without cracking. However, excessive single-pass downward pressure or an overall high rolling volume may cause cracks in the rolled plate at an early or late stage, while too low a rolling deformation may limit the formation of interface composites.

[0021] As a preferred solution, the annealing conditions are: atmospheric atmosphere, temperature of 623~773 K, and time of 3~30 min. The present invention controls the diffusion rate of elements in the Fe-Co-M based martensitic alloy by strictly controlling the annealing temperature and time, thereby regulating the formation of the interface complex, and ultimately obtaining an ideal high-density interface complex with effectively stable low-angle grain boundaries. If the annealing temperature is too low or the time is too short, it will lead to insufficient element diffusion, affecting the effect of regulating the interface stability and thus affecting the strengthening effect. If the annealing temperature is too high or the annealing time is too long, it will cause the element diffusion rate to be too fast, resulting in the formation of coarse precipitate phases, thereby seriously deteriorating the plasticity. A further preferred annealing temperature is 673~773K, and a further preferred annealing time is 5~20 min.

[0022] As a preferred solution, during the homogenization and annealing processes, cooling is performed by water quenching, which has the characteristics of fast cooling rate and can stabilize the alloy phase after homogenization and annealing.

[0023] The present invention provides a method for preparing an interface composite-strengthened Fe-Co-M-based martensitic alloy, comprising the following steps:

[0024] 1) Ultrasonic washing of high-purity metal raw materials Fe, Co, and M with dilute hydrochloric acid to remove the surface oxide layer, then ultrasonic washing with ethanol 2-3 times to remove the dilute hydrochloric acid on the metal surface, drying the surface alcohol with a baking lamp, and accurately weighing; drying the non-metallic raw materials and crushing them to a particle size of -1 mm;

[0025] 2) All accurately weighed metal raw materials are placed in a vacuum arc melting furnace for melting. The dried raw materials are stacked in descending order of melting point in a water-cooled copper crucible of the vacuum arc melting furnace, with the non-metallic Q at the bottom, Fe and Co in the middle, and Mo, which has the highest melting point, at the top. Vacuum melting is achieved in an arc melting furnace under the following conditions: an arc current of 370-390 A, a melting time of 90-120 seconds per melting, and after cooling for 4-5 minutes after each melting, the sample is flipped and the next melting is repeated 5-8 times. The melted samples are subjected to vacuum suction casting and copper mold cooling techniques to obtain rectangular columnar alloy specimens.

[0026] 3) Homogenize the vacuum-sealed cylindrical sample in a muffle furnace. The homogenization conditions are: in an inert gas atmosphere, at a temperature of 1373-1523 K, for 1.5-2 hours. After homogenization, quench with water.

[0027] 4) After homogenization, the surface oxide layer of the sample was removed with 80 molybdenum sandpaper, and then pre-deformed. The cold deformation treatment conditions were: rolling at room temperature, the total rolling deformation was 70-90%, the pass deformation was 3-5%, and the roller speed was 4-8 r / min.

[0028] 5) The pre-deformed sample is subjected to a low-temperature short-time annealing treatment to obtain the obtained product; the annealing conditions are: atmospheric atmosphere, temperature of 623~773 K, time of 3~30 min, and water quenching after annealing.

[0029] Compared with similar processes in the prior art, it has the following obvious advantages:

[0030] (1) The preparation method of the interface composite strengthened Fe-Co-M based martensitic alloy provided by the present invention has the advantages of simple process, short flow, safety, etc., and is conducive to industrial production.

[0031] (2) The main elements of the interface composite strengthened Fe-Co-M based martensitic alloy provided by the present invention are ordinary pure metal raw materials, and the production cost is relatively low.

[0032] (3) The interface complex-reinforced Fe-Co-M-based martensitic alloy provided by the present invention can effectively utilize the intrinsic characteristics of high-density dislocations and low-angle grain boundaries in the martensitic structure. The interface complex is introduced through a simple annealing process to improve the mechanical stability of the low-angle grain boundaries, thereby enhancing their resistance to dislocation migration and improving the alloy strength. This strategy greatly utilizes the characteristics of the microstructure in the martensitic structure, significantly increasing the number density of dislocation pinning points in the structure while retaining the dislocation permeability of the low-angle grain boundaries. Ultimately, it achieves a breakthrough in the strength of the martensitic structure alloy and excellent comprehensive performance of tensile ductility within the applicable range.

[0033] (4) The interface composite reinforced Fe-Co-M based martensitic alloy provided by the present invention can achieve a yield strength of 3.05 GPa while maintaining tensile plasticity within a usable range. BRIEF DESCRIPTION OF THE DRAWINGS

[0034] Figure 1 The alloys (Fe 49 Co 40 Mo 11 ) 99.6 B 0.3 C 0.1 Synchrotron radiation diffraction patterns under cold rolling (11MoBC-CR) and low temperature annealing (11MoBC-CRA) treatment conditions.

[0035] Figure 2 In comparative example 2 of the present invention, the alloy (Fe 49 Co 40 Mo 11 ) 99.6 B 0.3 C 0.1 SEM microstructure of (11MoBC-CR) after cold rolling.

[0036] Figure 3 The alloy used in Example 1 of the present invention (Fe 49 Co 40 Mo 11 ) 99.6 B 0.3 C 0.1 SEM microstructure of 11MoBC-CRA after cold rolling and annealing at 723K for 20min.

[0037] Figure 4 The alloy Fe used in Comparative Example 2 of the present invention 49 Co 40 Mo 11 APT microstructure morphology of (11Mo-CR) after cold rolling.

[0038] Figure 5 The alloy used in Example 1 of the present invention (Fe 49 Co 40 Mo 11 ) 99.6 B 0.3 C 0.1 APT microstructure of the interface composite after cold rolling and annealing at 723K for 20min (11MoBC-CRA).

[0039] Figure 6The alloys used in Example 1 and Comparative Example 2 (Fe 49 Co 40 Mo 11 ) 99.6 B 0.3 C 0.1 Mechanical properties under cold rolling (11MoBC-CR) and low temperature annealing (11MoBC-CRA) treatment conditions.

[0040] Figure 7 When the alloying element Mo (9 at.%) is added in Example 2 of the present invention, that is, (Fe 51 Co 40 Mo9) 99.6 B 0.3 C 0.1 -SEM microstructure of the alloy after cold rolling and annealing at 723K for 20 min (9MoBC-CRA).

[0041] Figure 8 When the alloying element Mo (9 at.%) is added in Example 2 of the present invention, that is, (Fe 51 Co 40 V9) 99.6 B 0.3 C 0.1 -Mechanical properties of the treated condition after cold rolling and annealing at 723K for 20 min (9MoBC-CRA).

[0042] Figure 9 When the alloying element V (10 at.%) is added in Example 3 of the present invention, that is, (Fe 50 Co 40 V 10 ) 99.6 B 0.3 C 0.1 -SEM microstructure of the alloy after cold rolling and annealing at 723K for 20 min (10VBC-CRA).

[0043] Figure 10 When the alloying element V (10 at.%) is added in Example 3 of the present invention, that is, (Fe 50 Co 40 V 10 ) 99.6 B 0.3 C 0.1 Mechanical properties of steel sheets annealed at 723K for 20 min (10VBC-CRA) after cold rolling.

[0044] Figure 11 When the alloying element TiVMoW (10 at.%) is added in Example 4 of the present invention, that is, (Fe 50 Co 40(TiVMoW) 10 ) 99.6 B 0.3 C 0.1 SEM microstructure of the alloy (10(TiVMoW)BC-CRA) annealed at 723K for 20 min after cold rolling.

[0045] Figure 12 When the alloying element TiVMoW (10 at.%) is added in Example 4 of the present invention, that is, (Fe 50 Co 40 (TiVMoW) 10 ) 99.6 B 0.3 C 0.1 Mechanical properties of the alloy (10(TiVMoW)BC-CRA) annealed at 723K for 20 min after cold rolling.

[0046] Figure 13 When there is no BC interstitial element in Comparative Example 1 of the present invention, that is, Fe 49 Co 40 Mo 11 Synchrotron radiation diffraction pattern of the alloy after cold rolling and annealing at 723K for 20 min (11Mo-CRA).

[0047] Figure 14 When there is no BC interstitial element in Comparative Example 1 of the present invention, that is, Fe 49 Co 40 Mo 11 SEM microstructure of the alloy after cold rolling and annealing at 723K for 20 min (11Mo-CRA).

[0048] Figure 15 When there is no BC interstitial element in Comparative Example 1 of the present invention, that is, Fe 49 Co 40 Mo 11 Mechanical properties of the alloy after cold rolling and annealing at 723K for 20 min (11Mo-CRA). DETAILED DESCRIPTION

[0049] To facilitate understanding of the present invention, the present invention will be described in more comprehensive and detailed form below in conjunction with the accompanying drawings and preferred embodiments. However, the scope of protection of the present invention is not limited to the following specific embodiments.

[0050] Unless otherwise defined, all technical terms used hereinafter have the same meanings as those generally understood by those skilled in the art. The technical terms used herein are only for the purpose of describing specific embodiments and are not intended to limit the scope of protection of the present invention.

[0051] Unless otherwise specified, various raw materials, reagents, instruments and equipment used in the present invention can be purchased from the market or prepared by existing methods.

[0052] Example 1

[0053] The present invention (Fe 49 Co 40 Mo 11 ) 99.6 B 0.3 C 0.1 (11MoBC-CRA alloy) martensitic alloy preparation and properties, including the following steps:

[0054] (1) High-purity (≥99.9%) Fe, Co, Mo, B, and C were used as alloy smelting raw materials. Each element was accurately weighed and proportioned: Fe: 26.681g, Co: 22.984g, Mo: 10.291g, B: 0.032g, and C: 0.012g, for a total weight of 60.000g. Each weighed metal raw material was then ultrasonically cleaned in alcohol and then placed on filter paper and dried with a heat lamp for later use. The non-metallic raw materials were dried and crushed to a particle size of -1mm.

[0055] (2) The dried raw materials are stacked in the water-cooled copper crucible of the vacuum arc melting furnace in descending order of melting point, with B and C placed at the bottom, Fe and Co placed in the middle, and Mo, which has the highest melting point, placed at the top. Then close the furnace door and evacuate the furnace chamber to 1×10 -3 Pa, and then fill the furnace with half atmospheric pressure. Then evacuate the furnace again to 1×10 -3 Pa, then fill the furnace with half an atmospheric pressure, ignite the tungsten electrode head, and start melting. Before melting the target alloy, first move the ignited tungsten electrode head to the titanium ingot and wait for it to melt for 30 seconds. The arc current for melting the Ti ingot is 130 A. The purpose of this step is to absorb free oxygen, nitrogen and other gases in the furnace. Finally, move the tungsten electrode head to the target alloy, gradually increase the current density to 390 A, and after all elements of the target alloy are melted, move the tungsten electrode head to the center of the ingot and stay for 120 seconds. After the time is up, gradually reduce the current, and then turn off the arc current. Wait for 5 minutes for the alloy to cool and then turn it over. Repeat this melting 8 times, and melt 4 times in the middle. After the alloy cools, evacuate the furnace chamber to 1×10 -3 After the alloy is fully melted and uniformly melted, the ingot alloy is taken out.

[0056] (3) The ingot alloy melted in step (2) was ground with a grinding wheel to remove the surface oxide layer, and then it was placed in alcohol and cleaned twice with ultrasonic vibration. It was then blown dry with a hair dryer and placed in a cleaned vacuum melting furnace. The copper mold was screwed into the mold and the alloy was suction-casted. Repeat the process twice to evacuate the furnace chamber to 1×10 -3 Pa, and then filled with 0.5 atmospheres of argon gas. After the Ti ingot was melted and held for 30 s, the tungsten electrode was moved to the target alloy ingot. The ingot was preheated for two cycles at 300 A, and then the current was increased to 400 A until the alloy was completely melted and flowed into the copper mold. The arc current was turned off to obtain a 10×10×80 mm martensitic alloy cuboid sample.

[0057] (4) Place the sample obtained by suction casting in step (3) into a glass tube with a groove at the waist and insert a glass stopper. Install the glass tube on a tube sealing machine and evacuate to 2×10 -3 After the sample reaches 100 Pa, it is filled with argon to half an atmosphere of pressure. This process is repeated at least five times. The glass tube and the glass stopper are then welded together securely using an oxyhydrogen generator. After the tube cools, it is placed in a muffle furnace for a homogenization treatment at 1473 K for two hours to obtain a uniform alloy. Following homogenization, the tube containing the alloy sample is placed in a metal bucket for water cooling. During this process, the tube is ensured to break in the bucket to ensure the sample is water-cooled.

[0058] (5) After homogenization in step (4), the surface oxide layer of the sample was removed using 80 molybdenum sandpaper or a sand tray, and then the sample was rolled at room temperature. The total rolling deformation was 85%, the rolling deformation per pass was 4%, and the roller speed was 6 r / min.

[0059] (6) The sample after rolling in step (5) is directly placed in a muffle furnace for annealing at 723K for 20 min under water cooling conditions. Finally, the target sample is obtained.

[0060] The product of this embodiment was tested by tissue observation method. 49 Co 40 Mo 11 ) 99.6 B 0.3 C 0.1 (11MoBC-CRA alloy) The phase composition of martensitic alloy after annealing is as follows Figure 1 As shown in the figure, the main phase composition is BCC martensite structure, with a small amount of austenite and boride phase. Figure 3As shown, the grains contain high-density dislocations and lath martensite interfaces, as well as a small amount of discretely distributed boride particles. This embodiment forms a Mo-BC interface complex at this annealing temperature, i.e., there is a co-segregation of the metal element Mo and the interstitial atoms B and C at the low-angle interface ( Figure 5 ). After testing (such as Figure 6 ), the yield strength of this material is: 3.05 GPa, and the tensile plasticity is: 5.1%.

[0061] Example 2

[0062] The present invention (Fe 51 Co 40 Mo9) 99.6 B 0.3 C 0.1 (9MoBC-CRA alloy) martensitic alloy preparation and properties, including the following steps:

[0063] (1) High-purity (≥99.9%) Fe, Co, Mo, B, and C were used as alloy smelting raw materials. Each element was accurately weighed and proportioned: Fe: 28.137g, Co: 23.288g, Mo: 8.531g, B: 0.032g, and C: 0.012g, for a total weight of 60.000g. Each weighed metal raw material was then ultrasonically cleaned in alcohol and then placed on filter paper and dried with a heat lamp for later use. The non-metallic raw materials were dried and crushed to a particle size of -1mm.

[0064] (2) The dried raw materials are stacked in the water-cooled copper crucible of the vacuum arc melting furnace in descending order of melting point, with B and C placed at the bottom, Fe and Co placed in the middle, and Mo, which has the highest melting point, placed at the top. Then close the furnace door and evacuate the furnace chamber to 1×10 -3 Pa, and then fill the furnace with half atmospheric pressure. Then evacuate the furnace again to 1×10 -3 Pa, then fill the furnace with half an atmospheric pressure, ignite the tungsten electrode head, and start melting. Before melting the target alloy, first move the ignited tungsten electrode head to the titanium ingot and stay there for 30s. The arc current for melting the Ti ingot is 130 A. The purpose of this step is to absorb free oxygen, nitrogen and other gases in the furnace. Finally, move the tungsten electrode head to the target alloy, gradually increase the current density to 390 A, and after all elements of the target alloy are melted, move the tungsten electrode head to the center of the ingot and stay there for 120 seconds. After the time is up, gradually reduce the current, and then turn off the arc current. Wait for 5 minutes for the alloy to cool and then turn it over. Repeat this melting 8 times, and melt 4 times in the middle. After the alloy cools down, evacuate the furnace chamber to 1×10 -3 After the alloy is fully melted and uniformly melted, the ingot alloy is taken out.

[0065] Steps (3) to (5) are consistent with those in Example 1.

[0066] (6) The sample after rolling in step (5) is directly placed in a muffle furnace for annealing at 723K for 20 min under water cooling conditions. Finally, the target sample is obtained.

[0067] The product of this embodiment was tested by tissue observation method. 51 Co 40 Mo9) 99.6 B 0.3 C 0.1 (9MoBC-CRA alloy) microstructure Figure 7 As shown in FIG, the grains contain high-density dislocations and lath martensite interfaces, as well as a small amount of discretely distributed carbide particles. In this embodiment, a Mo-BC interface complex with a relatively low Mo content is formed at this annealing temperature, i.e., there is a co-segregation of the metal element Mo and the interstitial atoms B and C at the small-angle interface. After testing (such as Figure 8 ), the yield strength of this material is: 2.690 GPa, and the tensile plasticity is: 3.6%.

[0068] Example 3

[0069] The present invention (Fe 50 Co 40 V 10 ) 99.6 B 0.3 C 0.1 (10VBC-CRA alloy) martensitic alloy preparation and properties, including the following steps:

[0070] (1) High-purity (≥99.9%) Fe, Co, V, B, and C were used as alloy smelting raw materials. Each element was accurately weighed and proportioned: Fe: 29.582g, Co: 24.974g, V: 5.397g, B: 0.034g, and C: 0.013g, for a total weight of 60.000g. Each weighed metal raw material was then ultrasonically cleaned in alcohol and then placed on filter paper and dried with a heat lamp for later use. The non-metallic raw materials were dried and crushed to a particle size of -1mm.

[0071] (2) The dried raw materials are stacked in the water-cooled copper crucible of the vacuum arc melting furnace in descending order of melting point, with B and C placed at the bottom, Fe and Co placed in the middle, and V, which has the highest melting point, placed at the top. Then close the furnace door and evacuate the furnace chamber to 1×10 -3 Pa, and then fill the furnace with half atmospheric pressure. Then evacuate the furnace again to 1×10 -3Pa, then fill the furnace with half an atmospheric pressure, ignite the tungsten electrode head, and start melting. Before melting the target alloy, first move the ignited tungsten electrode head to the titanium ingot and stay there for 30s. The arc current for melting the Ti ingot is 130 A. The purpose of this step is to absorb free oxygen, nitrogen and other gases in the furnace. Finally, move the tungsten electrode head to the target alloy, gradually increase the current density to 390 A, and after all elements of the target alloy are melted, move the tungsten electrode head to the center of the ingot and stay there for 120 seconds. After the time is up, gradually reduce the current, and then turn off the arc current. Wait for 5 minutes for the alloy to cool and then turn it over. Repeat this melting 8 times, and melt 4 times in the middle. After the alloy cools down, evacuate the furnace chamber to 1×10 -3 After the alloy is fully melted and uniformly melted, the ingot alloy is taken out.

[0072] Steps (3) to (5) are consistent with those in Example 1.

[0073] (6) The sample after rolling in step (5) is directly placed in a muffle furnace for annealing at 723K for 20 min under water cooling conditions. Finally, the target sample is obtained.

[0074] The product of this embodiment was tested by tissue observation method. 50 Co 40 V 10 ) 99.6 B 0.3 C 0.1 (10VBC-CRA alloy) microstructure Figure 9 As shown, the grains contain high-density dislocations and lath martensite interfaces, as well as a small amount of discretely distributed carbide particles. This embodiment forms a VBC interface complex at this annealing temperature, i.e., there is a co-segregation of the metal element V and the interstitial atoms B and C at the low-angle interface. Figure 10 ), the yield strength of this material is: 2.00 GPa, and the tensile plasticity is: 5.0%.

[0075] Example 4

[0076] The present invention (Fe 50 Co 40 (TiVMoW) 10 ) 99.6 B 0.3 C 0.1 The preparation and properties of the martensitic alloy (10(TiVMoW)BC-CRA alloy) include the following steps:

[0077] (1) High-purity (≥99.9%) Fe, Co, Ti, V, Mo, W, B, and C were used as alloy smelting raw materials. Each element was accurately weighed and proportioned: Fe: 27.462g, Co: 23.185g, Ti: 1.177g, V: 1.253g, Mo: 2.359g, W: 4.520g, B: 0.032g, and C: 0.012g, with a total weight of 60.000g. Each weighed metal raw material was then ultrasonically cleaned in alcohol and then placed on filter paper and dried with a heat lamp for later use. The non-metallic raw materials were dried and crushed to a particle size of -1mm.

[0078] (2) The dried raw materials are stacked in the water-cooled copper crucible of the vacuum arc melting furnace in descending order of melting point, with W powder, B and C placed at the bottom, Fe, Co and Ti placed in the middle, and V and Mo with the highest melting points placed at the top. Then close the furnace door and evacuate the furnace chamber to 1×10 -3 Pa, and then fill the furnace with half atmospheric pressure. Then evacuate the furnace again to 1×10 -3 Pa, then fill the furnace with half an atmospheric pressure, ignite the tungsten electrode head, and start melting. Before melting the target alloy, first move the ignited tungsten electrode head to the titanium ingot and stay there for 30s. The arc current for melting the Ti ingot is 130 A. The purpose of this step is to absorb free oxygen, nitrogen and other gases in the furnace. Finally, move the tungsten electrode head to the target alloy, gradually increase the current density to 390 A, and after all elements of the target alloy are melted, move the tungsten electrode head to the center of the ingot and stay there for 120 seconds. After the time is up, gradually reduce the current, and then turn off the arc current. Wait for 5 minutes for the alloy to cool and then turn it over. Repeat this melting 8 times, and melt 4 times in the middle. After the alloy cools down, evacuate the furnace chamber to 1×10 -3 After the alloy is fully melted and uniformly melted, the ingot alloy is taken out.

[0079] Steps (3) to (5) are consistent with those in Example 1.

[0080] (6) The sample after rolling in step (5) is directly placed in a muffle furnace for annealing at 723K for 20 min under water cooling conditions. Finally, the target sample is obtained.

[0081] The product of this embodiment was tested by tissue observation method. 50 Co 40 (TiVMoW) 10 ) 99.6 B 0.3 C 0.1 (10(TiVMoW)BC alloy) was tested (such as Figure 11), the grains contain high-density dislocations and lath martensite interfaces, and a small amount of discretely distributed carbide particles. In this embodiment, a Co-Ti-V-Mo-WBC interface complex is formed at this annealing temperature, that is, there is a co-segregation of seven elements, namely, metal elements Co, Ti, V, Mo, W, and interstitial atoms B and C, at the small-angle interface. After testing (such as Figure 12 ), its yield strength: 2.43 GPa, tensile plasticity: 4.7%.

[0082] Example 5 Alloy composition: (Fe 53 Co 40 Mo7) 99.6 B 0.3 C 0.1 .

[0083] The preparation method is completely based on Example 1.

[0084] The resulting alloy properties include yield strength of 2.48 GPa and fracture ductility of 3.43%.

[0085] Comparative Example 1

[0086] Fe of the present invention 49 Co 40 Mo 11 The preparation and properties of the martensitic alloy (11Mo-CRA alloy) include the following steps:

[0087] (1) High-purity (≥99.9%) Fe, Co, and Mo were used as alloying raw materials. Each element was accurately weighed and proportioned: 26.700g Fe, 23.001g Co, and 10.299g Mo, for a total weight of 60.000g. Each weighed metal raw material was then ultrasonically cleaned in alcohol and then placed on filter paper and dried using a heat lamp. The non-metallic raw materials were dried and crushed to a particle size of -1mm.

[0088] (2) The dried raw materials were stacked in a water-cooled copper crucible of a vacuum arc melting furnace in descending order of melting point, with Co at the bottom, Fe in the middle, and Mo having the highest melting point at the top; the remaining conditions and steps were the same as those in Example 1.

[0089] The product of this embodiment was tested by tissue observation method. 49 Co 40 Mo 11 (11Mo-CRA alloy) The martensitic alloy after annealing has a body-centered cubic (BCC) single phase, as determined by synchrotron radiation. Figure 13 Its organization is as shown. Figure 14 As shown in Figure 2, the grains contain high density dislocations and lath martensite interfaces. Figure 15 ), the yield strength of the material is 2.93GPa, but the tensile plasticity is only 1%. It can be seen that when the alloy does not contain B and C interstitial elements, under the same cold rolling and annealing temperature treatment conditions, the Fe 49 Co 40 Mo 11 The tensile ductility of the 11Mo-CRA alloy is severely deteriorated, significantly lower than 5% of that in Example 1. This is mainly because in this composition, only the Mo element segregates at the small-angle interface of the martensite, while no interstitial atoms segregate at the interface, severely reducing the interfacial cohesion and causing catastrophic early failure of the martensite alloy.

[0090] Comparative Example 2

[0091] The present invention (Fe 49 Co 40 Mo 11 ) 99.6 B 0.3 C 0.1 The preparation and properties of the martensitic alloy (11MoBC-CR alloy) include the following steps:

[0092] Steps (1) to (5) are the same as those in Example 1, and step (6) is omitted. Finally, the target sample is obtained.

[0093] The product of this embodiment was tested by tissue observation method. 49 Co 40 Mo 11 ) 99.6 B 0.3 C 0.1 (11MoBC-CR alloy) martensitic alloy phase composition after cold rolling is as follows Figure 1 As shown in the figure, it is mainly BCC phase structure with a small amount of austenite and boride phase. Figure 2 As shown, it is a cold-worked deformation state. The grains contain high-density dislocations and lath martensite interfaces, but there is no element segregation at the small-angle grain boundaries, and all elements are evenly distributed ( Figure 4 ). After testing (such as Figure 6 ), the yield strength of this material is 1.94 GPa, and the tensile plasticity is 6.5%. Compared to Example 1, when a non-interface composite is formed, the yield strength of the alloy material is significantly lower than that of Comparative Example 1. This is because the low-angle grain boundaries of martensite have limited dislocation barrier capacity, failing to fully contribute to strength, resulting in poor strengthening effect.

[0094] Comparative Example 3

[0095] The only difference between Comparative Example 3 and Example 1 is that the percentage of Mo in the M element is 15 at.%, and the remaining steps and conditions are the same. When proceeding to step (5), the alloy material cannot be rolled due to its excessive brittleness.

Claims

1. An interface complex-strengthened Fe-Co-M-based martensitic alloy, characterized in that: The element percentage composition expression is: Fe a Co b M c Q d , wherein, 40%≤a≤60%, 30%≤b≤50%, 0<c≤14%, 0<d≤2%, a+b+c+d=100; M includes at least one of Mo, V, Ti, W, Nb, Ta, Cr, Mn, Zr, Hf, Si, P, Ga, In, Sn, Pb, Ge, As, Sb, Te, Cu, Zn, Au, Ag, Pt, Pd, Cd, Ru, K, Mg, Y, and Ca; Q includes at least one of B, C, and N; its microphase comprises a high-density dislocation structure and a low-angle grain boundary structure modified by a complex composed of metal-interstitial atoms.

2. The method for preparing an interface composite strengthened Fe-Co-M based martensitic alloy according to claim 1, characterized in that: The metal raw materials M, Fe and Co are pre-treated for surface cleaning, and the non-metallic raw materials Q are pre-treated for crushing. The pre-treated metal raw materials and non-metallic raw materials are added into a melting furnace, and vacuum melting, casting, homogenization treatment, cold deformation treatment and annealing treatment are carried out in sequence to obtain the product.

3. The method for preparing an interface composite strengthened Fe-Co-M based martensitic alloy according to claim 2, characterized in that: The metal raw material is ultrasonically washed with alcohol; The non-metallic raw materials are crushed to a particle size of -1 mm.

4. The method for preparing an interface composite strengthened Fe-Co-M based martensitic alloy according to claim 2, characterized in that: During the vacuum smelting process, the metal raw materials are stacked in the smelting furnace from top to bottom in the order of melting points, and the non-metallic raw materials are placed at the bottom of the smelting furnace.

5. The method for preparing an interface composite strengthened Fe-Co-M based martensitic alloy according to claim 2 or 4, characterized in that: The vacuum melting is achieved by an arc melting furnace. The arc melting conditions are: arc current of 370-390 A, each melting time of 90-120 s, cooling for 4-5 minutes after each melting, turning the sample over, and then performing the next melting, and repeating the melting 5-8 times.

6. The method for preparing an interface composite strengthened Fe-Co-M based martensitic alloy according to claim 2, characterized in that: The homogenization treatment conditions are: in an inert gas atmosphere, at a temperature of 1373-1523 K, and for 1.5-2 h.

7. The method for preparing an interface composite strengthened Fe-Co-M based martensitic alloy according to claim 2, characterized in that: The cold deformation treatment conditions are: rolling at room temperature, a total rolling deformation of 70-90%, a pass deformation of 3-5%, and a roller running speed of 4-8 r / min.

8. The method for preparing an interface composite strengthened Fe-Co-M based martensitic alloy according to claim 2, characterized in that: The annealing treatment conditions are: atmospheric atmosphere, temperature of 623-773 K, and time of 3-30 min.

9. The method for preparing an interface composite strengthened Fe-Co-M based martensitic alloy according to claim 2, 6 or 8, characterized in that: During the homogenization treatment and the annealing treatment, cooling is performed by water quenching.

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